Low thermal expansion aluminum titanate - zirconium tin titanate ceramics

ABSTRACT

Disclosed herein is a ceramic body comprising at least one phase comprising a pseudobrookite-type crystal structure and at least one phase comprising zirconium tin titanate. Also disclosed are porous ceramic honeycomb structures comprising a ceramic body comprising at least one phase comprising a pseudobrookite-type crystal structure and at least one phase comprising zirconium tin titanate and methods of preparing a ceramic body comprising at least one phase comprising a pseudobrookite-type crystal structure and at least one phase comprising zirconium tin titanate.

CROSS-REFERENCE TO RELATED APPLICATION

This application is a divisional of and claims the benefit of priorityfrom U.S. patent application Ser. No. 13/920,493 filed on Jun. 18, 2013,the content of which is relied upon and incorporated herein by referencein its entirety.

TECHNICAL FIELD

The present disclosure relates to a ceramic body comprising at least onephase comprising a pseudobrookite-type crystal structure and at leastone phase comprising zirconium tin titanate. Also disclosed herein areporous ceramic honeycomb structures comprising a ceramic body comprisingat least one phase comprising a pseudobrookite-type crystal structureand at least one phase comprising zirconium tin titanate. Furtherdisclosed herein are methods of preparing a ceramic body comprising thesteps of providing a batch composition comprising at least one zirconiumsource, at least one tin source, at least one titanium source, at leastone aluminum source, and at least one magnesium source, and firing thebatch composition under conditions suitable to form a ceramic bodycomprising at least one phase comprising a pseudobrookite-type crystalstructure and at least one phase comprising zirconium tin titanate.

BACKGROUND

The after-treatment of exhaust gas from internal combustion engines mayrequire the use of filters and catalysts supported on high-surface areasubstrates. In the case of diesel engines and some gasoline directinjection engines, a catalyzed filter for the removal of carbon sootparticles may be used. The filters and catalyst supports in theseapplications should be refractory, thermal shock resistant, stable undera range of pO₂ conditions, non-reactive with the catalyst system, andshould offer minimal resistance to the exhaust gas flow. Porous ceramicflow-through honeycomb substrates and wall-flow honeycomb filters may,for example, be used in these applications.

SUMMARY

An exemplary embodiment of the disclosure provides a ceramic bodycomprising at least one phase comprising a pseudobrookite-type crystalstructure and at least one phase comprising zirconium tin titanate.

An exemplary embodiment of the disclosure also provides a porous ceramichoneycomb structure comprising a ceramic body comprising at least onephase comprising a pseudobrookite-type crystal structure and at leastone phase comprising zirconium tin titanate.

An exemplary embodiment of the disclosure also provides a method forpreparing a ceramic body, said method comprising the steps of providinga batch composition comprising at least one zirconium source, at leastone tin source, at least one titanium source, at least one aluminumsource, and at least one magnesium source, and firing the batchcomposition under conditions suitable to form a ceramic body comprisingat least one phase comprising a pseudobrookite-type crystal structureand at least one phase comprising zirconium tin titanate.

Both the foregoing general summary and the following detaileddescription are exemplary only and are not restrictive of thedisclosure. Further features and variations may be provided in additionto those set forth in the description. For instance, the disclosuredescribes various combinations and subcombinations of the featuresdisclosed in the detailed description. In addition, it will be notedthat where steps are disclosed, the steps need not be performed in thatorder unless explicitly stated.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a ternary composition diagram in mole percentages of ZrO₂,SnO₂, and TiO₂ components in which the region A-B-C-D-E-F-A enclosescompositions of zirconium tin titanate phases coexisting with Al₂TiO₅based pseudobrookite phases in certain embodiments of the ceramic bodiesdisclosed herein.

FIG. 2 shows a ternary composition diagram in mole percentages of ZrO₂,SnO₂, and TiO₂ components showing the approximate extent of solidsolution (ss) of the zirconium tin titanate based phase, “(Zr,Ti,Sn)₂O₄(ss),” at 1500-1600° C. The extent of solid solubility in the zirconiumoxide phase, “ZrO₂ (ss),” and the tin oxide and titanium oxide (rutile)phase, “SnO₂ (ss)” and “(Sn,Ti)O₂ (ss),” are only schematic and notmeasured. Also shown are symbols denoting specific compositions ofsingle-phase zirconium tin titanate compounds (filled circles) andcompositions that yield predominantly single-phase zirconium tintitanate compounds with minor amounts of zirconium oxide solid solution,“ZrO₂ (ss),” or tin oxide solid solution, “SnO₂ (ss),” (filled squares).The mean linear lattice CTE of the respective zirconium tin titanatecompounds from room temperature to 1000° C., as determined byhigh-temperature x-ray diffractometry, are shown next to the compositionsymbols. Dashed curves within the zirconium tin titanate phase fieldindicate compositions of constant CTE with values noted by arrows.

FIG. 3 is a plot of the measured four-point modulus of rupture (MOR)versus the calculated “S” parameter for pseudobrookite and zirconium tintitanate ceramics with not more than 2% rutile and not more than 15%corundum (filled circles); pseudobrookite and zirconium tin titanateceramics with more than 2% rutile (circles with vertical lines);pseudobrookite and zirconium tin titanate ceramics with more than 15%corundum (circles with horizontal lines); pseudobrookite and zirconiumtin titanate ceramics with more than 2% rutile and more than 15%corundum (circle with diagonal lines); pseudobrookite and tin-freezirconium titanate ceramics (open squares); Al₂TiO₅ based ceramics withno magnesium and no zirconium titanate phase (triangles with horizontallines); and pseudobrookite ceramics comprising 93.4 mole % Al₂TiO₅ and6.6 mole % MgTi₂O₅ with no zirconium titanate phase (open triangles).Lines indicate values of MOR/S=0.85, 1.00, and 1.15.

FIG. 4 shows a scanning electron micrograph of a polished cross sectionof Example 16 showing the homogeneously distributed zirconium tintitanate phase (white), the magnesium-containing Al₂TiO₅ basedpseudobrookite phase (gray), and the well-interconnected porosity(black).

FIG. 5 is a ternary composition diagram in mole percentages of ZrO₂,SnO₂, and TiO₂ components showing the compositions of zirconium tintitanate phases coexisting with an Al₂TiO₅ based pseudobrookite phaseand a corundum phase at 1400° C. (open circles) and at 1500° C. (filledcircles). The lines I-J, K-L, and M-N indicate the approximate minimumtitanium content in exemplary zirconium tin titanate phases that maycoexist in stable equilibrium with an Al₂TiO₅ based pseudobrookite phaseat 1400° C., 1500° C., and 1600° C., respectively.

FIG. 6 is a plot of the Young's elastic modulus versus temperatureduring heating (shown as circles) and cooling (shown as squares) forinventive Example 17, showing the measured elastic modulus at roomtemperature before heating (E_(RT)), measured elastic modulus atapproximately 1100° C. during cooling (E_(c, 1100)), and the estimatedvalue of the room-temperature elastic modulus for the sample in theabsence of microcracking (shown as a triangle, E°_(RT)).

FIG. 7 is a plot of the Young's elastic modulus versus temperatureduring heating (shown as circles) and cooling (shown as squares) forcomparative Example 12, showing the measured elastic modulus at roomtemperature before heating (E_(RT)), measured elastic modulus atapproximately 1100° C. during cooling (E_(c, 1100)), and estimated valueof the room-temperature elastic modulus for the sample in the absence ofmicrocracking (shown as a triangle, E°_(RT)).

FIG. 8 is a plot of the Young's elastic modulus versus temperatureduring heating (shown as circles) and cooling (shown as squares) forcomparative Example 10, showing the measured elastic modulus at roomtemperature before heating (E_(RT)), measured elastic modulus atapproximately 1100° C. during cooling (E_(c, 1100)), and estimated valueof the room-temperature elastic modulus for the sample in the absence ofmicrocracking (shown as a triangle, E°_(RT)).

DESCRIPTION OF EXEMPLARY EMBODIMENTS

Since the introduction of substrates in the 1970s and filters in the1980s, there has been an increasing trend toward thinner walls in bothsubstrates and filters to reduce pressure drop. Furthermore, there isalso a demand for higher porosity substrates to reduce thermal mass tofacilitate earlier activation of the catalyst during engine cold startsand higher porosity in filters to further reduce pressure drop and toaccommodate larger amounts of catalysts within the filter wall. Thinnerwalls and higher porosity both may weaken the structure of the honeycombmaterial; therefore, it is desirable that the inherent flexural strengthof the skeletal ceramic material comprising the walls be as high aspossible.

Also, in the case of diesel particulate filters (DPFs), a highvolumetric heat capacity may be desired in order to minimize thetemperatures that may be reached during the periodic in situ combustionof the accumulated soot, a process known as “regeneration.” Since higherporosity and thinner walls reduce the mass of the DPF, the solidmaterial comprising the ceramic should have a high heat capacity perunit volume. To address this need, aluminum titanate based DPFs havebeen introduced.

Some of the new catalysts with improved performance that are availablehave been found to exhibit undesirable chemical interactions withexisting ceramic filters and substrates. For example, potassium-basednitric oxide (NO_(x)) adsorbers may be rapidly deactivated when appliedto silicate-containing ceramics, such as cordierite or aluminumtitanate+feldspar. Also, the activity of copper-containing chabazitezeolite catalysts has been found to be reduced by even small amounts ofsodium that may occur as an impurity in many cordierite and aluminumtitanate+feldspar ceramics.

Moreover, certain catalysts may also have detrimental effects on theceramic support. For example, it has been found that copper may diffuseor leach from a chabazite catalyst and promote decomposition of aluminumtitanate into corundum and rutile, with a subsequent loss in thermalshock resistance of the aluminum titanate+feldspar filter. Ceramicscomprised of cordierite+mullite+aluminum titanate are being developed toprovide a more durable filter in the presence of copper; however,relative to previous aluminum titanate+feldspar filters, these materialsmay have a lower volumetric heat capacity, lower melting point, and/orrequire the use of expensive rare earth oxide additions as sinteringaids. Furthermore, such ceramics may contain a glassy grain-boundaryphase, and it is proposed that reduction or elimination of agrain-boundary silicate phase may be beneficial in reducing diffusionrates of undesirable cations from the ceramic to the catalyst or fromthe catalyst to the ceramic.

In an effort to respond to the current trends in reduced wall thickness,higher porosity, higher thermal mass for DPF applications, and reducedchemical interaction with new catalyst systems, disclosed herein arenovel ceramic materials based upon aluminum titanate+zirconium tintitanate that, in at least certain embodiments, may exhibit superiorstrength, higher volumetric heat capacity, an absence of a grainboundary silicate phase, and/or that may not require the use ofexpensive sintering aids.

In at least certain embodiments disclosed herein, the presence of tin inthe zirconium tin titanate phase may be beneficial for providingimproved strength to the ceramic body at a low coefficient of thermalexpansion (CTE) for a given porosity and pore size. Specifically, thepresence of tin in the zirconium tin titanate phase of the ceramicbodies disclosed herein that also contain an aluminum titanatepseudobrookite phase may impart a greater flexural strength, at a given% porosity, median pore diameter, and CTE, than ceramics comprised of analuminum titanate pseudobrookite phase and a tin-absent zirconiumtitanate phase. Without wishing to be bound by theory, it is believedthat this may be the result of the lower CTE of zirconium tin titanaterelative to tin-free zirconium titanate, such that the ceramic bodiesdisclosed herein require less microcracking to achieve a given low CTE.

Moreover, the ceramic bodies disclosed herein may not require the use ofsintering additives to achieve the desired properties. For example, inat least certain embodiments, the ceramic bodies disclosed herein may besintered at temperatures ranging from about 1400-1500° C. without theneed for sintering additives, making them highly refractory. Inparticular, the expensive rare earth oxide sintering additives used inmany other aluminum titanate based ceramics may not be required incertain embodiments disclosed herein.

In certain embodiments, the ceramic bodies disclosed herein do not havea grain-boundary silicate phase. The absence of a grain-boundarysilicate phase may render the ceramic bodies useful in applications inwhich an alkali, such as potassium, is present, for example either inash deposited from exhaust produced by combustion of a biodiesel fuel,or as a NO_(x)-adsorbing catalyst system. The absence of agrain-boundary silicate phase may also inhibit the exchange of otherchemical elements between the ceramic body and any applied catalyst,thereby improving the lifetime of both the ceramic body and thecatalyst.

In certain embodiments disclosed herein, the at least one phasecomprising a pseudobrookite-type crystal structure contains MgTi₂O₅.Embodiments disclosed herein wherein the at least one phase comprising apseudobrookite-type crystal structure is predominantly Al₂TiO₅ with atleast several mole percent MgTi₂O₅ may exhibit excellenthigh-temperature stability in the presence of copper, making suchceramic bodies useful in applications in which a copper-containingcompound, such as a copper-containing chabazite zeolite, is applied tothe ceramic body as a catalyst. Such embodiments may have greater copperdurability than magnesium-free aluminum titanate based ceramics.

Zirconium titanate based compounds possess a volumetric heat capacity(4.30 J cm⁻³ K⁻¹ at 800K) that is higher than that of aluminum titanate(3.97 J cm⁻³ K⁻¹ at 800K), mullite (3.67 J cm⁻³ K⁻¹ at 800K), feldspar(2.99 J cm⁻³ K⁻¹ at 800K), and cordierite (2.78 J cm⁻³ K⁻¹ at 800K).Therefore, ceramic bodies disclosed herein that comprise at least onephase comprising a pseudobrookite-type crystal structure, wherein thatpseudobrookite phase is aluminum rich, and at least one phase comprisingzirconium tin titanate, are expected to have a higher volumetric heatcapacity than ceramics based upon aluminum titanate with a second phasecomprising at least one of mullite, feldspar, and cordierite. A highervolumetric heat capacity may be beneficial for application as a dieselparticulate filter, for example, wherein the greater thermal mass of thefilter reduces the temperatures experienced by the filter during asevere regeneration (combustion of accumulated carbon soot).

Disclosed herein are ceramic bodies comprising at least one phasecomprising a pseudobrookite-type crystal structure and at least onephase comprising zirconium tin titanate. In certain embodimentsdisclosed herein, the at least one phase comprising zirconium tintitanate has a crystal structure based upon, and isostructural with, thecompound ZrTiO₄ which, in turn, is based upon the structure of a-PbO₂.Also disclosed herein are porous ceramic honeycomb structures, forexample exhaust particulate filters such as diesel particulate filters,comprising a ceramic body comprising at least one phase comprising apseudobrookite-type crystal structure and at least one phase comprisingzirconium tin titanate. Exemplary porous ceramic honeycomb structuresmay include wall-flow, partial wall-flow, and flow-through honeycombstructures. In certain other exemplary embodiments, the ceramic body maybe non-porous.

Further disclosed herein are methods of preparing a ceramic bodycomprising the steps of providing a batch composition comprising atleast one zirconium source, at least one tin source, at least onetitanium source, at least one aluminum source, and at least onemagnesium source, and firing the batch composition under conditionssuitable to form a ceramic body comprising at least one phase comprisinga pseudobrookite-type crystal structure and at least one phasecomprising zirconium tin titanate. In certain embodiments, the batchcomposition may be fired at a temperature of at least about 1400° C.,such as at least about 1500° C. In certain embodiments, the batchcomposition may further comprise at least one pore forming agent. Thepore forming agents (pore formers) can include, for example, graphite,starch, nut shells, synthetic organic particles, or even combinationsthereof. The starch can include, for example, sago palm starch, greenmung bean starch, canna starch, corn starch, rice starch, pea starch, orpotato starch. The median particle diameter of the pore forming agent isselected according to the application of the ceramic honeycomb, and ispreferably between 1 and 60 microns, and is more preferably between 5and 40 microns.

Pseudobrookite-type crystal structure describes a phase having a crystalstructure that is the same as at least one of the mineralspseudobrookite (Fe₂TiO₅), karooite (MgTi₂O₅), and tialite (Al₂TiO₅),without restriction as to composition. The composition of the phasecomprising a pseudobrookite-type crystal structure may in certainembodiments be of the general formula(M²⁺)_(w)(M³⁺)_(x)(M⁴⁺)_(y)(M⁵⁺)_(z)O₅, wherein M²⁺ is a divalent cationof an element such as magnesium, cobalt, nickel, and copper; M³⁺ is atrivalent cation of an element such as aluminum, iron, gallium,scandium, and titanium; M⁴⁺ is a tetravalent cation of an element suchas titanium, zirconium, tin, silicon, and germanium; M⁵⁺ is apentavalent cation of an element such as niobium and tantalum;(w+x+y+z)=3; and 2w+3x+4y+5z=10. It is understood that, as used herein,all instances of w, x, y and z are expressed as molar quantities.

In certain embodiments, the pseudobrookite-type phase is of thecomposition (Mg)_(w1)(Al)_(x1)(Fe)_(x2)(Ti)_(y1)(Zr)_(y2)(Sn)_(y3)O₅ inwhich 2(w1)+3(x1+x2)+4(y1+y2+y3)=10 and wherein 0.0≦w1≦0.50,0.95≦x1≦2.01, 0.0≦x2≦0.05, 0.70<y1≦1.5, 0.0<y2≦0.10, and 0.0<y3≦0.20. Incertain embodiments, 0.0≦w1≦0.15, 1.70≦x1≦2.01, 0.0≦x2≦0.01,0.83<y1≦1.09, 0.01<y2≦0.05, and 0.05<y3≦0.12. In certain otherembodiments disclosed herein, the value of w1 is at least about 0.03, asmagnesium has been shown to be beneficial in slowing the rate at which aAl₂TiO₅-rich pseudobrookite phase decomposes below about 1280° C., forexample in the presence of copper ions, which may be present in certaincatalyst systems, such as those comprising a copper-containing chabazitezeolite catalyst, when the ceramic body disclosed herein is used as asubstrate or particulate filter in the after-treatment of exhaust froman internal combustion engine. In certain other embodiments, x2<0.001.It is understood that, as used herein, all instances of w1, x1, x2, y1,y2 and y3 are expressed as molar quantities.

The zirconium tin titanate has a crystal structure that may be describedas based upon, and isostructural with, the compound ZrTiO₄. In certainembodiments, the ZrTiO₄ in the ZrO₂—TiO₂ system may undergo at least onephase transformation associated with cation ordering in the crystallattice at lower temperatures; the zirconium tin titanate phasedisclosed herein is not restricted as to the extent of cation ordering.In certain embodiments, however, the crystal structure of the zirconiumtin titanate phase may be similar to that of the high-temperature,cation-disordered form of ZrTiO₄.

The composition of the at least one phase comprising zirconium tintitanate may, in certain embodiments, be of the general formula(M′³⁺)_(p1)(Zr)_(c1)(Ti)_(q2)(Sn)_(q3)(Hf)_(q4)(M″⁵⁺)_(r)O_(4-s) whereinM′³⁺ is a trivalent cation of an element such as aluminum, gallium, andiron; M″⁵⁺ is a pentavalent cation of an element such as tantalum andniobium; (p1+q1+q2+q3+q4+r)=2; q3≧0.05; and 0≦s≦0.05.

In certain other embodiments, the zirconium tin titanate phase is of thecomposition (Al)_(p1)(Zr)_(q1)(Ti)_(q2)(Sn)_(q3)O_(4-0.5(p1)), wherein(p1+q1+q2+q3)=2; 0<p1≦0.08; and q3≧0.05. It is understood that, as usedherein, all instances of p1, q1, q2, q3, q4, r and s are expressed asmolar quantities. In other embodiments disclosed herein, the value of q2is not more than about 1.00, such as about 0.90, about 0.80, about 0.70,and about 0.60. In certain other embodiments disclosed herein, the valueof q3 may be at least about 0.10, such as at least about 0.20, at leastabout 0.30, at least about 0.40, at least about 0.50, and at least about0.60. In other embodiments, the value of q1/(q1+q2+q3) is between about0.36 and about 0.60, the value of q2/(q1+q2+q3) is between about 0.23and about 0.50, and the value of q3/(q1+q2+q3) is between about 0.05 andabout 0.33. The compositional region defined by these limits isrepresented, for example, in FIG. 1 by the polygon A-B-C-D-E-F-A,wherein the proportions of ZrO₂, SnO₂, and TiO₂ are given in molepercentages. In at least certain embodiments, the zirconium tin titanatephase contains at least about 5 mole % SnO₂.

The ceramic bodies disclosed herein may, in certain embodiments, have aporosity greater than about 45%, such as, for example, greater thanabout 50%, greater than about 60%, greater than about 65%, or rangingfrom about 50% to about 65%. In certain embodiments disclosed herein,the ceramic body has a median pore diameter ranging from about 11 μm toabout 15 μm. In other embodiments disclosed herein, the ceramic body hasa narrow pore size distribution with values of (d₉₀−d₁₀)/d₅₀,(d₉₀−d₅₀)/d₅₀, (d₅₀−d₁₀)/d₅₀ all less than about 0.30. In otherexemplary embodiments disclosed herein, the ceramic body may have aCTE_(RT-1000° C.) value ranging from about −15×10⁻⁷° C.⁻¹ to about+30×10 ⁻⁷° C.⁻¹, making the ceramic bodies disclosed herein useful aswall-flow and partial wall-flow particulate filters for after-treatmentof gasoline or diesel engine exhaust. In certain embodiments, theparticle sizes of the raw materials and the amounts and particle sizesof organic pore formers could be adjusted to produce porous ceramicbodies for use as flow-through substrates, such as flow-throughcatalytic converter substrates.

Certain embodiments disclosed herein may include applications requiringa high thermal shock resistance. For such embodiments, the ceramic bodydisclosed herein may have a low CTE and a high value ofMOR/[(E)(CTE_(500-1000° C.))], where MOR and E (Young's elastic modulus)are measured at room temperature. The elastic modulus may be measured bya sonic resonance method. The quantity MOR/[(E)(CTE_(500-1000° C.))] isproportional to the thermal shock resistance of the ceramic body.Specifically, the value of MOR/[(E)(CTE_(500-1000° C.))] provides anestimate of the temperature difference, ΔT, between the hotter interiorand the cooler outer surface of the ceramic body that the body canwithstand before fracturing when the temperature of the outer surface ofthe ceramic body is equal to 500° C. Therefore, high ratios of MOR/E andlow values of CTE_(500-1000° C.) may provide for a high thermal shockresistance. For certain embodiments, the value of CTE_(500-1000° C.) forthe ceramic bodies disclosed herein may be generally proportional to thevalue of CTE_(RT-1000° C.). Therefore, a low value of CTE_(RT-1000° C.)may also achieve a high thermal shock resistance. Accordingly, incertain embodiments disclosed herein, the value of CTE_(RT-1000° C.) isless than about 30×10⁻⁷° C.⁻¹, less than about 25×10⁻⁷° C.⁻¹, less thanabout 20×10⁻⁷° C.⁻¹, less than about 15×10⁻⁷° C.⁻¹, or even less thanabout 10×10⁻⁷° C.⁻¹. In certain embodiments, the value ofMOR/[(E)(CTE_(500-1000° C.))] is at least about 250° C., at least about300° C., at least about 400° C., at least about 500° C., at least about600° C., at least about 800° C., and even at least about 1000° C. Avalue of MOR/[(E)(CTE_(500-1000° C.))] equal to 400° C., for example,indicates that, when the outer surface of the ceramic body is at 500°C., the interior of the ceramic body may be heated to 900° C. (500°C.+ΔT=500° C.+400° C.) before fracture from thermally induced stressesoccurs.

The presence of tin in the zirconium tin titanate phase, which may, incertain exemplary embodiments, be in combination with low amounts oftitanium, has been found to be beneficial for enabling improved valuesof MOR for a given CTE, % porosity, and median pore size. This may beimportant in certain embodiments in which the porosity is high (such asgreater than about 50%), the value of d₅₀ is large (such as greater thanabout 12 μm), and the CTE is low (such as less than about 25×10⁻⁷°C.⁻¹), as low CTE may be associated with a high degree of microcracking,and high porosity, large pore size, and high microcracking all tendto/reduce strength. Specifically, in certain embodiments disclosedherein, the four-point flexural strength, MOR, of the ceramic, in unitsof pounds/inch² (psi), is greater than a value “S” where “S” is definedas in Equation 1 (EQ. 1) below:

S=2140−18.1(% porosity)−57.2(d ₅₀)+6.1(CTE_(RT-1000° C.))  EQ. 1:

wherein d₅₀ is the median pore diameter of the ceramic body in units ofmicrometers and % porosity is the volume percentage of porosity, both asmeasured by mercury porosimetry, and CTE_(RT-1000° C.) is the meancoefficient of thermal expansion between room temperature (approximately22° C.) and 1000° C. in units of 10⁻⁷° C.⁻¹ as measured by dilatometry.In certain embodiments, the ratio of the measured flexural strength tothe computed value of “S”, (MOR/S), is at least about 1.05, at leastabout 1.10, at least about 1.15, at least about 1.20, and even at leastabout 1.25. A high value of MOR/S means that the ceramic body desirablyhas an especially high strength at a given set of values for % porosity,median pore diameter, and coefficient of thermal expansion.

Certain embodiments disclosed herein comprising at least one phasecomprising zirconium tin titanate have a lower CTE at a given degree ofmicrocracking relative to embodiments where tin is absent from thezirconium titanate phase. Likewise, certain embodiments disclosed hereinmay not require as much microcracking to provide a similar CTE as anembodiment wherein tin is absent from the zirconium titanate phase.

Microcracking may be the result of tensile stresses that develop betweenadjacent crystals in the ceramic body during cooling when the crystalshave different CTEs from one another or when the crystals exhibitanisotropic CTEs along their different crystallographic directions.Although two ceramic bodies having approximately the same % porosity,median pore diameter, and CTE, will have approximately the samecalculated “S” value, the embodiments disclosed herein that contain azirconium tin titanate phase have been found to possess lessmicrocracking and a higher MOR and higher MOR/S ratio. Therefore,embodiments disclosed herein may have a greater strength for a given CTEthan embodiments where tin is absent from the zirconium titanate phase.Additionally, embodiments disclosed herein may have a lower CTE at thesame level of microcracking as embodiments where tin is absent from thezirconium titanate phase, resulting in a greater thermal shockresistance for a given strength.

During heating of a ceramic body, the microcracks, which originallyformed during the cooling of the ceramic body after firing, maygradually re-close. This process of re-closing may become more rapidabove about 800° C. By about 1200° C., most of the microcracks have beeneliminated by closure and annealing, resulting in a stiffening of theceramic body and an increase in elastic modulus. During the initialstages of cooling from about 1200° C., the elastic modulus remains highbecause stresses between grains in the ceramic body are too low to causemicrocracking. Below about 1000° C., however, the stresses between thegrains increase to the point where microcracking may occur, and theelastic modulus may decrease with further cooling.

If a tangent line is drawn to the elastic modulus values during theinitial stages of the cooling curve, such as at about 1100° C., then theextrapolation of this tangent line to room temperature yields the valueof the elastic modulus of the ceramic body at room temperature in ahypothetical non-microcracked state (E°_(RT)). See, for example, FIG. 6.The slope of the tangent line thus represents the change in the elasticmodulus of a non-microcracked version of the ceramic body withtemperature change, ΔE°/ΔT. This slope is negative because, duringcooling, the stiffening of the atomic bonds within the crystal grainscauses the elastic modulus of the crystals to increase. The value ofE°_(RT) and the slope of the tangent may also be affected by % porosity.However, the ratio of the slope to the value of E°_(RT) is a constant.

The ratio of the room-temperature elastic modulus of a hypotheticalceramic body in the non-microcracked state, E°_(RT), to the value of theroom-temperature elastic modulus of an actual microcracked ceramic body,E_(RT), is proportional to the amount of microcracking within theceramic body, as more microcracking produces a lower value of E_(RT)relative to the value of E°_(RT).

In certain embodiments disclosed herein wherein the MOR is greater thanthe value of S, the ceramic body may comprise a microstructure whereinthe raw materials have undergone complete or nearly complete reaction toform the desired pseudobrookite and zirconium tin titanate phases. Thisis because complete reaction may be associated with a high degree ofsintering in which the individual crystals comprising the ceramic bodyare well bonded to one another, which is to say that a high proportionof the crystal surfaces are bonded to the surfaces of adjacent crystals,thereby maximizing the connectivity and contiguity of the solid phase ata given percentage of porosity in the ceramic body. Accordingly, theceramic body should, in certain embodiments, be heated to a sufficientlyhigh temperature and held at that temperature for a sufficient time suchthat, after firing, the ceramic body contains a minimum amount of rutile(crystalline TiO₂ phase) and a minimum amount of corundum.

It has been found that the presence of corundum in certain embodimentsdisclosed herein may be the result of a reaction between thepseudobrookite phase and the zirconium tin titanate phase during firing.It has been discovered that for a zirconium tin titanate phase to be inequilibrium with (i.e., coexist with) an Al₂TiO₅-based pseudobrookitephase during firing, the zirconium tin titanate phase should contain acertain minimum amount of titanium. If the zirconium tin titanate phasedoes not contain this minimum amount of titanium, it may react with theAl₂TiO₅-rich pseudobrookite phase so as to incorporate the titanium fromthe Al₂TiO₅-rich pseudobrookite phase into the crystal structure of thezirconium tin titanate phase, thereby also forming a free corundum(aluminum oxide) phase as one of the reaction products.

To avoid this reaction, it has been found that the amount of titanium inthe zirconium tin titanate phase should be, according to certainembodiments disclosed herein, at least about 28-32 cation % when thematerial is fired at about 1500° C., and, in certain embodiments, atleast about 35-38 cation % when the material is fired at about 1400° C.This is equivalent to about 28-32 mole % TiO₂ component in the zirconiumtin titanate phase when the material is fired at about 1500° C., or atleast about 35-38 mole % TiO₂ component in the zirconium tin titanatephase when the material is fired at about 1400° C. Any attempt to form aceramic comprising an Al₂TiO₅-rich pseudobrookite phase and a zirconiumtin titanate phase having less than the minimum required titaniumcontent may, instead, result in a ceramic comprising an Al₂TiO₅-richpseudobrookite phase with a corundum phase and a zirconium tin titanatephase having a titanium content equal to the minimum amount required forequilibrium with the Al₂TiO₅-rich pseudobrookite phase at the firingtemperature.

For example, in certain embodiments, the ceramic body may contain lessthan about 2.0 weight percent of rutile and less than about 2.0 weightpercent corundum. Higher amounts of rutile in combination with at leastabout 2.0 weight percent corundum may indicate incomplete formation ofthe pseudobrookite phase. In certain embodiments, incomplete formationof the pseudobrookite phase may be detrimental to the strength (MOR) ofthe ceramic body.

In order to achieve the best combination of high strength (high MOR/S)and low CTE, in certain embodiments disclosed herein, the fired ceramicbody contains no more than about 15 wt % corundum. It is believed thatlarger amounts of corundum may contribute to an increased CTE of theceramic due to the intrinsically high coefficient of thermal expansionof corundum, which is approximately 85×10⁻⁷° C.⁻¹. It is speculated thatthe high CTE of corundum may also contribute to increased microcracking,which would contribute to a lower MOR, as discussed herein. Therefore,in certain embodiments, the ceramic body disclosed herein may notcontain more than about 10 wt % corundum, and even not more than about 5wt % corundum or not more than about 2 wt % corundum.

Furthermore, in order to provide a ceramic body comprising anAl₂TiO₅-based pseudobrookite phase and a zirconium tin titanate phase inwhich the zirconium tin titanate phase has the lowest possible meanlattice coefficient of thermal expansion, in certain embodiments, onemay select a bulk composition and a firing temperature that will allowfor the stable coexistence of the Al₂TiO₅-based pseudobrookite phasewith a zirconium tin titanate phase having the lowest possible titaniumcontent without formation of more than about 15 wt % corundum byreaction with the pseudobrookite phase during firing.

Unless otherwise indicated, all numbers used in the specification andclaims are to be understood as being modified in all instances by theterm “about,” whether or not so stated. It should also be understoodthat the precise numerical values used in the specification and claimsform additional embodiments of the disclosure. Efforts have been made toensure the accuracy of the numerical values disclosed in the Examples.Any measured numerical value, however, can inherently contain certainerrors resulting from the standard deviation found in its respectivemeasuring technique.

As used herein the use of “the,” “a,” or “an” means “at least one,” andshould not be limited to “only one” unless explicitly indicated to thecontrary.

It is to be understood that both the foregoing general description andthe detailed description are exemplary and explanatory only and are notintended to be restrictive.

The accompanying drawings, which are incorporated in and constitute apart of this specification, are not intended to be restrictive, butrather illustrate embodiments of the disclosure.

Other embodiments will be apparent to those skilled in the art fromconsideration of the specification and practice of the disclosure.

EXAMPLES

The following examples are not intended to be limiting of thedisclosure.

Example Unit Cell Dimensions

Before the ceramic bodies discussed below in Examples 1-30 wereprepared, a series of compositions in the ZrO₂—TiO₂—SnO₂ system weresynthesized at temperatures of 1500° C. and 1600° C. to define the limitof ZrTiO₄ solid solution in the ternary system at these temperatures andto quantify the changes in the unit cell dimensions and latticecoefficients of thermal expansion as a function of the composition.Samples were prepared from mixtures of powders of the end-member oxides,which were blended with methyl cellulose and plasticized by the additionof water in a stainless steel muller. These materials were then extrudedas 8 mm diameter rod using a ram extruder, dried, and fired in anelectric furnace by heating at 50° C./h, holding 10 hours at maximumtemperature, and cooling at 500° C./h. The phases in the fired materialand their weight percentages were determined by powder x-raydiffractometry using a Rietveld refinement of the data.

The crystal lattice axial CTE values (i.e., the CTE values along thedirections of the “a,” “b,” and “c” unit cell dimensions) weresubsequently determined by high-temperature x-ray diffractometry forthose compositions which yielded single-phase, or nearly single-phase,(Zr,Ti,Sn)₂O₄ materials when fired at 1500-1600° C. These compositionsare listed in Table 1, along with their room-temperature unit celldimensions and their lattice CTE values. The mean lattice CTE value isthe average of the three crystal lattice axial CTE values. Theapproximate extent of solid solution of the (Zr,Ti,Sn)₂O₄ phase at 1600°C. is depicted in FIG. 2. Also shown in FIG. 2 are the locations of thespecific compositions in Table 1, as well as their mean lattice CTEvalues. A series of dashed contour lines denote the approximate locationof compositions having the same CTE values of 40, 50, 60, 70, and80×10⁻⁷° C.⁻¹. The data in Table 1 and FIG. 2 demonstrate the reductionin the mean lattice CTE of (Zr,Ti,Sn)₂O₄ with decreasing titaniumcontent and increasing tin content. At the highest ratios of Sn/Ti, themean lattice CTE is further reduced with increasing zirconium content.

The unit cell dimensions listed in Table 1 were fit by least-squaresmultiple linear regression analysis to yield the following threeequations, in which “Zr,” “Sn,” and “Ti” are the number of zirconium,tin, and titanium atoms, respectively, in a four-oxygen (two cation)formula unit:

“a” unit cell dimension (Å)=4.56705+0.239524(Zr)+0.061806(Sn) 98.5%R²  EQ. 2:

“b” unit cell dimension(Å)=5.07079+0.21806(Zr)+0.707332(Sn)−0.321695(Sn)²+0.13948(Ti)² 99.8%R²  EQ. 3:

“c” unit cell dimension (Å)=5.25636−0.0489722(Zr)−0.182123(Ti) 99.8%R²  EQ. 4:

Table 1 below shows the values of the a, b, and c unit cell dimensionsat room temperature, and mean lattice coefficients of thermal expansionof the a, b, and c unit cell dimension between room temperature and1000° C., and the average of the three lattice CTE values, for zirconiumtitanate and various zirconium tin titanate compounds. The compositionsare indicated by the number of atoms per four-oxygen formula unit, andCTE values are in units of 10⁻⁷° C.⁻¹.

TABLE 1 Formula Zr/4ox Ti/4ox Sn/4ox a (Å) b (Å) c (Å) CTE (a) CTE (b)CTE (c) CTE (avg) ZrTiO₄ 1.00 1.00 0.00 4.8059 5.4287 5.0274 69.3 103.387.1 86.6 Zr_(0.75)TiSn_(0.25)O₄ 0.75 1.00 0.25 4.7564 5.5321 5.039285.3 58.3 76.9 73.5 Zr_(0.75)Ti_(0.75)Sn_(0.5)O₄ 0.75 0.75 0.50 4.77875.5826 5.0810 88.2 35.9 47.3 57.1 ZrTi_(0.7)Sn_(0.3)O₄ 1.00 0.70 0.304.8290 5.5408 5.0792 85.5 37.9 63.7 62.4 Zr_(1.20)Ti_(0.60)Sn_(0.20)O₄1.20 0.60 0.20 4.8745 5.5076 5.0835 81.9 38.3 63.4 61.2ZrTi_(0.4)Sn_(0.6)O₄ 1.00 0.40 0.60 4.8453 5.6211 5.1351 101.4 2.4 44.149.3 Zr_(1.20)Ti_(0.40)Sn_(0.40)O₄ 1.20 0.40 0.40 4.8750 5.5922 5.1257113.4 −23.4 40.3 43.4 ZrTi_(0.3)Sn_(0.7)O₄ 1.00 0.30 0.70 4.8500 5.64085.1531 104.1 −1.5 38.4 47.0 ZrTi_(0.2)Sn_(0.8)O₄ 1.00 0.20 0.80 4.86145.6547 5.1694 108.2 −7.0 34.2 45.1 Zr_(1.20)Ti_(0.20)Sn_(0.60)O₄ 1.200.20 0.60 4.8823 5.6420 5.1646 117.1 −26.5 29.8 40.1

Examples 1-29

Examples of embodiments disclosed herein, based upon aluminumtitanate+zirconium tin titanate or magnesium aluminum titanate+zirconiumtin titanate, and comparative examples, based upon aluminumtitanate±tin-free zirconium titanate or magnesium aluminumtitanate±tin-free zirconium titanate, are presented in Tables 2 to 6below.

End member oxide powders were used with the exception that a spinel(MgAl₂O₄) powder was employed to supply the magnesium. Median particlediameters, as determined by laser diffraction, are listed for each ofthe inorganic powders. In addition to the inorganic constituents, 12parts by weight of a graphite powder (49 μm median particle diameter)and 22 parts by weight of a cross-linked pea starch were added to 100parts by weight of the inorganic powders to serve as pore formers. Thiswas done in an effort to yield porosity levels of approximately 50-60%in the fired ware, which is a range that may be appropriate for use ofthe ceramic bodies as a diesel particulate filter, although the ceramicbodies disclosed herein are not limited by % porosity or by application.

To the combination of inorganic and pore-former powders were added 4.5parts by weight of methyl cellulose and 1.0 parts by weight of tall oil,and sufficient water was added to the mixture in a stainless steelmuller to provide a plasticized batch. Each batch was subsequentlyloaded into the chamber of a ram extruder, de-aired by pulling a vacuumon the chamber, and pushed through a die to form 60 cm lengths of 8 mmdiameter rod. The rod was dried at 85° C. in open-ended glass tubes, cutinto 3-inch and 6-inch lengths, placed in alumina trays, and fired inair in an electric furnace at 50° C./h to either 1400 or 1500° C., heldfor 15 hours, and cooled at 500° C./h to room temperature.

The % porosity and pore size distribution of the fired ware weredetermined by mercury porosimetry. Tables 2-6 list the values of thed₁₀, d₅₀, and d₉₀ pore sizes (in micrometers), which are the porediameters at 10%, 50% and 90% of the cumulative pore size distributioncurve based upon pore volume, such that d₉₀ is the pore diameter atwhich 10% of the pore volume of the sample has been intruded by mercury,d₅₀ is the pore diameter at which 50% of the pore volume of the samplehas been intruded by mercury, and d₁₀ is the pore diameter at which 90%of the pore volume of the sample has been intruded by mercury. Thus,d₁₀<d₅₀<d₉₀.

Also provided in Tables 2 to 6 are the values of (d₉₀−d₁₀)/d₅₀,(d₉₀−d₅₀)/d₅₀, and (d₅₀−d₁₀)/d₅₀, which are normalized measures of theoverall breadth of the pore size distribution, the breadth of the coarseend of the distribution, and breadth of the fine end of thedistribution, respectively. Lower values of these metrics correspond tonarrower pore size distributions, and a narrow pore size distribution,especially in the fine half of the distribution curve, may be beneficialfor providing low pressure drop in the soot-loaded state, for examplewhen the ceramic body is used as a diesel particulate filter. A narrowpore size distribution, especially in the coarse half of thedistribution curve, may also be beneficial for improved flexuralstrength, as low values of (d₉₀−d₅₀)/d₅₀ may imply an absence of largepores that can serve as strength-limiting flaws.

The thermal expansion of the ceramic body along the length of a 2-inchrod was measured from room temperature to 1000° C. using a push-roddilatometer, and values of the mean (“secant”) coefficients of thermalexpansion between room temperature and 1000° C., and between 500° C. and1000° C., both during heating, are listed in Tables 2 to 6.

Modulus of rupture was measured at room temperature on 3-inch long rods'using the four-point method with a 0.75-inch load span and a 2.0-inchsupport span.

Young's elastic modulus values of certain examples were measured at roomtemperature using a sonic resonance technique. For selected examples,elastic modulus was further measured by a sonic resonance technique fromroom temperature up to 1200° C., and back to room temperature, atintervals of approximately 50° C.

The weight percentages of all crystalline phases in the fired ceramicswere determined by powder x-ray diffractometry and applying a Rietveldanalysis to the data, which also provided the unit cell parameters ofthe pseudobrookite and zirconium titanate based phases. Equations 2 to 4were used to estimate the number of atoms of Zr, Sn, and Ti in theformula units of the zirconium titanate and zirconium tin titanatephases based upon their unit cell parameters, using a least-squaresiterative procedure in which the values of “Zr,” “Sn,” and “Ti” wereadjusted until the unit cell dimensions predicted from the equationsexhibited the minimum deviation from the measured values.

Selected samples were also examined by scanning electron microscopy, andthe compositions of the pseudobrookite and zirconium titanate andzirconium tin titanate phases determined directly by electron probemicroanalysis.

For some examples, the thermal stability of the ceramic body in thepresence of copper (II) oxide was characterized. First, a quantity ofthe ceramic was pulverized for 30 seconds in a bench-top ring mill to amedian particle diameter of about 20 μm to about 40 μm. Three to fivegrams of pulverized ceramic were transferred to a 12 ml polypropylenecontainer, and 0.25 wt % of copper (II) oxide powder from Sigma Aldrichwas added. The container was placed on a SPEX Sample Prep 8000MMixer/Mill and shaken for 5 minutes to homogenize the powders. Themixture was transferred to an alumina crucible and heated in an electricfurnace at 120° C./h to 800° C., 300° C./h to 1100° C., held at 1100° C.for 2 hours, and cooled at 300° C./h to room temperature. The weightpercentages of pseudobrookite, corundum, rutile, and zirconium titanatebased phase after heat treatment were measured by x-ray diffractometryusing Rietveld refinement. The stability of the pseudobrookite phase inthe presence of the copper oxide was quantified by taking the ratio ofweight percent pseudobrookite phase after the copper test to the weightpercent of pseudobrookite phase in the original as-fired ceramic.

Examples of comparative materials to those of inventive ceramic bodiesare provided in Tables 2 and 3.

The comparative examples in Table 2 are porous ceramic bodies comprisedmainly of a pseudobrookite-type phase with small amounts of residualcorundum or rutile, but lacking a zirconium titanate based phase. Thepseudobrookite phase in Examples 1 and 2 is essentially pure Al₂TiO₅,and in Examples 3 and 4 the pseudobrookite phase is approximately 93.4mole % Al₂TiO₅ and 6.6 mole % MgTi₂O₅. The ratios of the MOR values ofthese comparative examples to their calculated “S” values were less than1.0, and their MOR values were plotted against their “S” values in FIG.3 (triangles).

The comparative examples in Table 3 represent ceramic bodies comprisedof a magnesium-free Al₂TiO₅ phase and a tin-free zirconium titanatephase (Examples 5 to 7) or a magnesium-containing Al₂TiO₅-based phaseand a tin-free zirconium titanate phase (Examples 8 to 12). Electronprobe microanalysis of the phases in Example 6 showed the presence of asmall amount of zirconium in the pseudobrookite phase and a small amountof aluminum in the zirconium titanate phase. The presence of the largerzirconium cation in the pseudobrookite phase results in the slightlylarger “a” and “b” unit cell dimensions of the pseudobrookite phaserelative to those of the Al₂TiO₅ in Example 2, wherein no zirconium ispresent. The even larger unit cell dimensions of the pseudobrookitephases in Examples 8 to 12 are the result of the presence of magnesiumin the crystal structure. The ratios of the MOR values to the computed“S” values for Examples 5 to 12 are all less than 1.0, ranging from 0.67to 0.88, and the MOR values are plotted against “S” values in FIG. 3(shown as open squares).

Examples of embodiments disclosed herein comprising a magnesium-freeAl₂TiO₅-based pseudobrookite phase and a zirconium tin titanate phase(Examples 13 and 14) and further embodiments comprising amagnesium-containing Al₂TiO₅-based pseudobrookite phase and a zirconiumtin titanate phase (Examples 15 to 23) for which the ratio MOR/S isgreater than 1.0 are provided in Tables 4 and 5.

The examples in Tables 4 and 5 encompass values of porosity (51 to 61%),median pore diameters (11 to 15 μm), (d₉₀−d₁₀)/d₅₀ (0.24 to 0.50),(d₅₀−d₁₀)/d₅₀ (0.17 to 0.43), and CTE (−15 to +29×10⁻⁷° C.⁻¹) which maybe appropriate for use as a diesel particulate filter. However,variation in any of these physical properties to produce a ceramic bodycomprising at least one phase comprising pseudobrookite and at least onephase comprising zirconium tin titanate for other applications could beachieved through changes in raw material particle size, raw materialmineralogy, and firing cycle, for example, without departing from thescope of the embodiments disclosed herein.

The values of MOR/S for the inventive examples of Tables 4 and 5 rangefrom 1.07 to 1.27, and MOR is plotted versus the computed “S” parameterin FIG. 3 (shown as solid-filled circles). Comparison with the data fromTables 2 and 3 shows that the presence of a zirconium tin titanium oxidephase in these inventive examples results in a greater flexural strengthfor a given % porosity, d₅₀, and CTE.

A scanning electron micrograph of a polished cross section of inventiveExample 16 is provided in FIG. 4, showing the uniform distribution ofthe zirconium tin titanate phase (brighter regions) and the highlyuniform porosity (black regions).

Examples 24 to 26 of Table 6 also represent inventive ceramic bodiescomprising predominantly a magnesium-containing Al₂TiO₅-richpseudobrookite phase and a zirconium tin titanate phase. However, thesubstantial amounts of corundum in these samples, comprising greater 15%by weight corundum, are associated with lower values of MOR/S, whichare, for example, less than 1.0. The MOR and S values of these examplesare plotted in FIG. 3 (shown as circles with horizontal or diagonallines), from which the lower values of strength for a given “S” valueare apparent. It has been found that high MOR/S values in the inventiveceramic bodies may require that the amount of corundum in the fired wareshould be, for example, less than about 15 wt %.

As discussed above, it has been found that the presence of corundum inthe inventive compositions, which ranges from about 2% to about 15% inthe ceramic bodies of Tables 4 and 5 and about 6% to about 22% in theceramic bodies of Table 6, is largely the result of a reaction betweenthe pseudobrookite phase and the zirconium tin titanate phase duringfiring.

To illustrate this point, the batch composition of Example 18 wascalculated to form a ceramic comprising 67 mole % of a pseudobrookitephase having the nominal compositionMg_(0.066)Al_(1.868)Zr_(0.05)Ti_(0.92)Sn_(0.1005) and 33 mole % of azirconium tin titanate phase having the nominal compositionZr_(1.17)Ti_(0.39)Sn_(0.39)Al_(0.05)O_(3.975). However, after firingExample 18 at 1500° C. for 15 hours, the composition of the zirconiumtin titanate phase was determined by electron probe microanalysis to beapproximately Zr_(0.96)Ti_(0.62)Sn_(0.38)Al_(0.04)O_(3.95) and that ofthe pseudobrookite-type phase was approximatelyMg_(0.108)Al_(1.824)Zr_(0.03)Ti_(0.95)Sn_(0.06)O₅. The sample alsocontained 12 weight % corundum, but no rutile. The abundance of corundumand absence of rutile indicates that the corundum is not the result ofincomplete reaction, but has formed as a stable reaction product betweenthe pseudobrookite and zirconium tin titanate phase. The titaniumcontent of the actual zirconium tin titanate phase is substantiallygreater than that of the targeted composition for that phase, and themagnesium content of the pseudobrookite-type phase is also higher thantargeted. These compositional differences between the actual andtargeted phase compositions are the result of extraction of the titaniumfrom the pseudobrookite-type phase and its dissolution into thezirconium tin titanate phase, and the expulsion of the excess alumina inthe pseudobrookite-type phase to form corundum, whereby the ratio ofMgTi₂O₅ component to Al₂TiO₅ component in the pseudobrookite-type phaseincreased, resulting in the higher magnesium content of thepseudobrookite-type phase.

FIG. 5 illustrates the compositions of zirconium tin titanate phases, asdetermined by either electron probe microanalysis or from the unit celldimensions measured by x-ray diffractometry, for some of the inventiveceramic examples from Tables 4 to 6 that also contain an Al₂TiO₅-richpseudobrookite phase and corundum. The compositions of the zirconium tintitanate phases from the samples fired at 1400° C. are denoted by opencircles, and those from samples fired at 1500° C. are denoted by filledcircles. The solid line I-J indicates the approximate lower limit oftitanium in zirconium tin titanate phases that can stably coexist withan Al₂TiO₅-based pseudobrookite phase at 1400° C., and the long-dashedline K-L indicates the approximate lower limit of titanium in zirconiumtin titanate phases that can stably coexist with an Al₂TiO₅-basedpseudobrookite phase at 1500° C. It is apparent that the range ofcompositions of the zirconium tin titanate phase that can coexist withan Al₂TiO₅-based pseudobrookite phase expands toward lower titaniumcontents with increasing firing temperature. The short-dashed line M-Nindicates the speculative lower limit of titanium in zirconium tintitanate phases that might stably coexist with an Al₂TiO₅-basedpseudobrookite phase fired at 1600° C. It will be noted that Examples 24to 26 in Table 6, fired at 1400° C., contain 18 to 22 wt % corundum andthe estimated titanium contents of the zirconium tin titanate phases are0.69 to 0.75 atoms per formula unit, whereas Examples 18 to 20 of Table5, which represent the same bulk compositions fired at 1500° C., containonly 9 to 15 wt % corundum and the zirconium tin titanate phases containonly 0.56 to 0.61 atoms of titanium per formula unit. This furtherillustrates the extension of the compositions of the zirconium tintitanate phase in equilibrium with an Al₂TiO₅-based pseudobrookite phaseto lower titanium contents with increasing firing temperature.

Accordingly and with reference to FIG. 5, in order to minimize theamount of corundum in the embodiments disclosed herein, the bulkcomposition of the ceramic body may be formulated to comprise only anAl₂TiO₅-based pseudobrookite phase and a zirconium tin titanate phasehaving a composition lying within the area bounded by G-H-I-J-G (andwith at least 2.5 mole % SnO₂ component) if the body is to be fired at1400° C., or within the area bounded by G-H-I-K-L-J-G (and with at least2.5 mole % SnO₂ component) if the body is to be fired at 1500° C., orwithin the area bounded by G-H-I-K-M-N-L-J-G (and with at least 2.5 mole% SnO₂ component) if the body is to be fired at 1600° C.

Examples 27 to 29 in Table 6 demonstrate that when the ceramic body isnot fully reacted such that there remains at least 2.0 wt % rutile andat least 2.0 wt % corundum in the fired ware, the value of MOR/S wasfound to be less than 1.0.

To show the effect of microcracking on both strength and CTE, the extentof microcracking in three examples (inventive Example 17 and comparativeExamples 10 and 12) was estimated from high-temperature elastic modulusmeasurements. In FIGS. 6-8, Young's elastic modulus is shown as afunction of temperature during heating to about 1200° C. and coolingback to room temperature. A value of −1.5×10⁻⁴° C.⁻¹ has been adoptedfor the value of (ΔE°/ΔT)/E°_(RT) in FIGS. 6-8, where the superscript“°” indicates the ceramic in a non-microcracked state. FIGS. 6, 7, and 8show plots of the Young's elastic modulus versus temperature duringheating (shown as circles) and cooling (shown as squares) for inventiveExample 17, comparative Example 12, and comparative Example 10,respectively. The plots illustrate the measured elastic modulus at roomtemperature before heating (E_(RT)), measured elastic modulus atapproximately 1100° C. during cooling (E_(c, 1100)), and estimated valueof the room-temperature elastic modulus for the sample in the absence ofmicrocracking (triangle, E°_(RT)).

Inventive Example 17 was compared with comparative Example 12. Bothsamples had similar values of CTE and % porosity. Example 17 contained26.3 weight percent zirconium tin titanate of approximate compositionZr_(0.95)Sn_(0.42)Ti_(0.63)O₄, while Example 12 contained 26.0 weightpercent of ZrTiO₄ (see Tables 3 and 4). The value of E°_(RT) for Example17 was determined to be 2.57×10⁶ psi, and the value of E_(RT) wasdetermined to be 2.93×10⁵ psi, giving a ratio E°_(RT)/E_(RT)=8.8. Thevalue of E°_(RT) for Example 12 was determined to be 1.63×10⁶ psi, andthe value of E_(RT) was determined to be 1.27×10⁵ psi, giving a ratioE°_(RT)/E_(RT)=12.8. The higher E°_(RT)/E_(RT) ratio of comparativeExample 12 demonstrates a higher level of microcracking in that sample.This greater amount of microcracking causes the MOR of comparativeExample 12, which is 230 psi, to be lower than that of inventive Example17, which is 393 psi. Also, the ratio of MOR/S for comparative Example12, 0.81, is lower than that of inventive Example 17, which is 1.07. Thehigher ratio of MOR/S of the inventive Example 17 indicates that, for agiven % porosity, median pore diameter, and CTE, the inventive exampleis stronger than the comparative example, and this is the result of thelower degree of microcracking in the inventive example coupled with thelower mean lattice CTE of the zirconium tin titanate phase in Example 17relative to that of the zirconium titanate phase in Example 12.

Inventive Example 17 was also compared with comparative Example 10. Bothexamples have approximately the same values of % porosity and MOR, butExample 10 has a higher CTE. The high-temperature elastic modulus datafor Example 10 is shown in FIG. 8. The value of E°_(RT) for Example 10has been determined to be 3.38×10⁶ psi, and the value of E_(RT) wasdetermined to be 3.70×10⁵ psi, giving a ratio E°_(RT)/E_(RT)=9.1.Therefore, Example 10 has only slightly more microcracking thaninventive Example 17. However, the CTE_(RT-1000° C.) of Example 10 is24.6×10⁻⁷° C.⁻¹, 50% higher than that of inventive Example 17. Thisdemonstrates that the embodiments disclosed herein comprising azirconium tin titanate phase have a lower CTE at a given level ofmicrocracking than the comparative bodies that lack tin in the zirconiumtitanate phase. The lower CTE of inventive Example 17 results in greaterpredicted thermal shock parameter MOR/(E·CTE_(500-1000° C.))=492° C.,compared to only 325° C. for comparative Example 10 at about the samelevel of microcracking. The lower CTE for a given level of microcrackingin the preferred embodiment of Example 17 is again reflected in thehigher value of MOR/S, 1.07, relative to that of comparative Example 10,which is 0.88.

TABLE 2 Comparative examples that do not contain a ZrTiO₄ phase ExampleNumber 1 2 3 4 Inorganic Raw Materials Al₂O₃ 12.0 μm  56.08 56.08 46.2846.28 ZrO₂ A 5.7 μm 0 0 0 0 ZrO₂ S 17.3 μm  0 0 0 0 SnO₂ 5.5 μm 0 0 0 0TiO₂ 0.5 μm 43.92 43.92 46.76 46.76 Spinel 6.5 μm 0 0 6.96 6.96 FiringConditions Hold Temperature (° C.) 1400 1500 1400 1500 Hold Time (hours)15 15 15 15 Fired Properties % Porosity 60.1 59.8 60.4 63.6 d₁₀ 9.8 11.010.3 12.0 d₅₀ 14.5 14.9 15.7 16.0 d₉₀ 15.8 16.6 16.9 17.5 (d₅₀ −d₁₀)/d₅₀ 0.32 0.27 0.35 0.25 (d₉₀ − d₅₀)/d₅₀ 0.09 0.11 0.07 0.09 (d₉₀ −d₁₀)/d₅₀ 0.41 0.37 0.42 0.34 CTE_(RT-1000° C.) (10⁻⁷° C.⁻¹) −9.1 −10.38.1 10.8 CTE_(500-1000° C.) (10⁻⁷° C.⁻¹) −12.6 −14.1 4.8 6.8 MOR (psi)136 114 188 134 “S” value (psi) 168 142 196 139 MOR/S ratio 0.81 0.810.96 0.96 Proportion and Composition of Phases wt % Pseudobrookite —97.1 95.5 98.7 wt % Corundum — 2.3 2.8 1.2 wt % Rutile — 0.6 1.7 0 Wt %(Zr,Ti,Sn)O₄ — 0 0 0 Pseudobrookite Lattice — 9.431 9.455 9.454Parameter a Pseudobrookite Lattice — 9.642 9.666 9.664 Parameter bPseudobrookite Lattice — 3.594 3.602 3.601 Parameter c Al per Psbformula unit — 2.007 — 1.873 from EPMA Ti per Psb formula unit — 0.993 —1.060 from EPMA Sn per Psb formula unit — 0.000 — 0.000 from EPMA Zr perPsb formula unit — 0.000 — 0.000 from EPMA Mg per Psb formula unit —0.000 — 0.066 from EPMA

TABLE 3 Comparative examples that contain a tin-free ZrTiO₄ phaseExample Number 5 6 7 8 9 10 11 12 Inorganic Raw Materials Al₂O₃ 12.0 μm 44.86 33.65 22.43 34.93 26.20 26.20 26.20 31.58 ZrO₂ A 5.7 μm 12.1324.27 36.40 12.13 0 24.27 24.27 0 ZrO₂ S 17.3 μm  0 0 0 0 24.27 0 022.46 SnO₂ 5.5 μm 0 0 0 0 0 0 0 0 TiO₂ 0.5 μm 43.00 42.08 41.16 45.8444.21 44.21 44.21 42.84 Spinel 6.5 μm 0 0 0 7.10 5.32 5.32 5.32 3.13Firing Conditions Hold Temperature (° C.) 1500 1500 1500 1500 1400 14001500 1500 Hold Time (hours) 15 15 15 15 15 15 15 15 Fired Properties %Porosity 51.9 48.8 44.6 52.4 60.5 58.5 50.2 57.7 d₁₀ 11.2 10.7 10.1 12.39.1 10.1 12.1 12.7 d₅₀ 13.8 12.9 12.0 15.1 13.6 13.4 14.2 15.9 d₉₀ 14.713.6 12.7 15.9 15.8 14.2 14.9 17.0 (d₅₀ − d₁₀)/d₅₀ 0.19 0.17 0.16 0.190.33 0.24 0.15 0.20 (d₉₀ − d₅₀)/d₅₀ 0.07 0.06 0.06 0.05 0.16 0.06 0.050.07 (d₉₀ − d₁₀)/d₅₀ 0.26 0.23 0.21 0.24 0.49 0.30 0.20 0.27CTE_(RT-1000° C.) (10⁻⁷° C.⁻¹) −11.3 −4.4 8.2 5.8 27.2 24.6 12.3 15.9CTE_(500-1000° C.) (10⁻⁷° C.⁻¹) −0.6 6.2 17.4 15.0 36.5 34.0 23.1 25.0MOR (psi) 275 432 543 249 288 408 396 230 “S” value (psi) 344 491 696363 431 464 494 284 MOR/S ratio 0.80 0.88 0.78 0.69 0.67 0.88 0.80 0.81Elastic Modulus at RT (psi) — — 4.82E+05 1.69E+05 2.18E+05 3.70E+052.94E+05 1.25E+05 MOR/(E*CTE_(500-1000° C.)) — — 647 980 362 325 584 725Proportion and Composition of Phases wt % Pseudobrookite 83.6 67.2 57.784.3 67.1 64.5 68.4 73.0 wt % Corundum 0.7 0.7 0 0.7 3.0 2.0 0.5 0.0 wt% Rutile 0.1 0.2 0 0.9 1.7 1.2 1.0 1.0 Wt % ZrTiO₄ 15.6 32.0 41.7 13.828.0 32.0 29.8 26.0 Pseudobrookite Lattice Parameter a 9.435 9.436 9.4349.470 9.470 9.470 9.469 9.459 Pseudobrookite Lattice Parameter b 9.6499.650 9.649 9.689 9.687 9.686 9.688 9.675 Pseudobrookite LatticeParameter c 3.592 3.592 3.591 3.602 3.602 3.601 3.601 3.598 ZrTiO₄Lattice Parameter a 4.798 4.805 4.807 4.805 4.802 4.803 4.807 4.822ZrTiO₄ Lattice Parameter b 5.412 5.417 5.411 5.416 5.408 5.411 5.4115.418 ZrTiO₄ Lattice Parameter c 5.025 5.028 5.028 5.028 5.025 5.0265.028 5.034 Approximate Zr per ZrTiO₄ from unit cell 0.99 1.01 1.02 1.011.00 1.00 1.02 1.07 Approximate Ti per ZrTiO₄ from unit cell 1.01 0.990.98 0.99 1.00 1.00 0.98 0.93 Approximate Sn per ZrTiO₄ from unit cell0.00 0.00 0.00 0.00 0.00 0.00 0.00 0.00 Al per Psb formula unit fromEPMA — 2.006 — — — — — — Ti per Psb formula unit from EPMA — 0.975 — — —— — — Sn per Psb formula unit from EPMA — 0.000 — — — — — — Zr per Psbformula unit from EPMA — 0.019 — — — — — — Mg per Psb formula unit fromEPMA — 0.000 — — — — — — Zr per ZTT formula unit from EPMA — 0.929 — — —— — — Ti per ZTT formula unit from EPMA — 1.025 — — — — — — Sn per ZTTformula unit from EPMA — 0.000 — — — — — — Al per ZTT formula unit fromEPMA — 0.046 — — — — — — Mg per ZTT formula unit from EPMA — 0.000 — — —— — — Oxygen per ZTT formula unit from EPMA — 3.977 — — — — — —

TABLE 4 Exemplary embodiments that contain a zirconium tin titanatephase Example Number 13 14 15 16 17 Inorganic Raw Materials Al₂O₃ 12.0μm  44.86 33.65 27.77 27.77 26.20 ZrO₂ A 5.7 μm 10.03 20.07 20.07 20.070 ZrO₂ S 17.3 μm  0 0 0 0 20.07 SnO₂ 5.5 μm 7.36 14.73 14.73 14.73 14.73TiO₂ 0.5 μm 37.74 31.56 33.26 33.26 33.69 Spinel 6.5 μm 0 0 4.18 4.185.32 Firing Conditions Hold Temperature (° C.) 1500 1500 1400 1500 1500Hold Time (hours) 15 15 15 15 15 Fired Properties % Porosity 55.3 53.860.6 51.3 57.3 d₁₀ 10.9 10.3 6.3 9.3 11.7 d₅₀ 13.9 12.5 11.0 11.8 14.6d₉₀ 15.1 13.4 11.7 12.6 15.6 (d₅₀ − d₁₀)/d₅₀ 0.22 0.17 0.43 0.21 0.20(d₉₀ − d₅₀)/d₅₀ 0.09 0.07 0.07 0.07 0.07 (d₉₀ − d₁₀)/d₅₀ 0.31 0.24 0.500.28 0.27 CTE_(RT-1000° C.) (10⁻⁷° C.⁻¹) −14.5 7.3 28.7 9.4 16.4CTE_(500-1000° C.) (10⁻⁷° C.⁻¹) 2.4 17.6 37.8 20.7 27.3 MOR (psi) 304631 697 701 393 “S” value (psi) 253 495 588 595 367 MOR/S ratio 1.201.27 1.19 1.18 1.07 Elastic Modulus at RT (psi) 2.22E+05 4.82E+055.79E+05 5.14E+05 2.93E+05 MOR/(E*CTE_(500-1000° C.)) * = estimated 5803743 319 660 492 Proportion and Composition of Phases wt % Pseudobrookite80.9 59.8 48.2 62.3 69.2 wt % Corundum 2.5 6.6 9.8 2.9 3.6 wt % Rutile0.4 1.0 0.5 0.2 1.0 Wt % (Zr,Ti,Sn)O₄ 16.2 35.6 41.6 34.5 26.3Pseudobrookite Lattice Parameter a 9.451 9.457 9.485 9.486 9.497Pseudobrookite Lattice Parameter b 9.671 9.680 9.709 9.713 9.726Pseudobrookite Lattice Parameter c 3.592 3.591 3.600 3.598 3.601(Zr,Ti,Sn)O₄ Lattice Parameter a 4.798 4.813 4.779 4.817 4.820(Zr,Ti,Sn)O₄ Lattice Parameter b 5.548 5.569 5.567 5.576 5.574(Zr,Ti,Sn)O₄ Lattice Parameter c 5.068 5.087 5.071 5.093 5.094Approximate Zr per ZrTiO₄ from unit cell 0.88 0.92 0.78 0.93 0.95Approximate Ti per ZrTiO₄ from unit cell 0.80 0.68 0.81 0.64 0.63Approximate Sn per ZrTiO₄ from unit cell 0.33 0.40 0.41 0.43 0.42 Al perPsb formula unit from EPMA — — — 1.880 — Ti per Psb formula unit fromEPMA — — — 0.960 — Sn per Psb formula unit from EPMA — — — 0.065 — Zrper Psb formula unit from EPMA — — — 0.028 — Mg per Psb formula unitfrom EPMA — — — 0.066 — Zr per ZTT formula unit from EPMA — — — 0.885 —Ti per ZTT formula unit from EPMA — — — 0.636 — Sn per ZTT formula unitfrom EPMA — — — 0.431 — Al per ZTT formula unit from EPMA — — — 0.048 —Mg per ZTT formula unit from EPMA — — — 0.000 — Oxygen per ZTT formulaunit from EPMA — — — 3.976 —

TABLE 5 Exemplary embodiments that contain a zirconium tin titanatephase, fired at 1500° C. for 15 hours Example Number 18 19 20 21 22 23Inorganic Raw Materials Al₂O₃ 12.0 μm  28.33 28.09 27.72 25.89 31.0031.31 ZrO₂ A 5.7 μm 0 0 0 0 0 0 ZrO₂ S 17.3 μm  24.98 20.97 24.44 19.4521.81 24.20 SnO₂ 5.5 μm 14.25 18.77 18.52 15.25 11.26 7.32 TiO₂ 0.5 μm28.64 28.40 25.60 33.00 32.96 34.18 Spinel 6.5 μm 3.81 3.78 3.73 6.412.97 3.00 Fired Properties % Porosity 59.1 59.5 56.3 57.7 59.2 58.7 d₁₀10.9 10.8 10.4 12.0 11.6 11.6 d₅₀ 14.1 13.9 13.4 14.9 14.9 15.2 d₉₀ 15.214.9 14.3 15.7 15.7 16.4 (d₅₀ − d₁₀)/d₅₀ 0.23 0.22 0.22 0.19 0.22 0.23(d₉₀ − d₅₀)/d₅₀ 0.08 0.07 0.07 0.05 0.05 0.08 (d₉₀ − d₁₀)/d₅₀ 0.31 0.290.29 0.24 0.28 0.31 CTE_(RT-1000°C) (10⁻⁷° C.⁻¹) 26.6 23.0 28.5 19.617.1 13.9 CTE_(500-1000°C) (10⁻⁷° C.⁻¹) 36.2 32.8 37.3 29.2 24.6 6.8 MOR(psi) 515 462 604 430 348 355 “S” value (psi) 425 409 529 365 319 293MOR/S ratio 1.21 1.13 1.14 1.18 1.09 1.21 Elastic Modulus at RT (psi)4.56E+05 4.45E+05 5.68E+05 3.01E+05 2.58E+05 2.45E+05MOR/(E*CTE_(500-1000° C.)) 312 316 286 489 550 2129 Proportion andComposition of Phases wt % Pseudobrookite 56.0 61.0 52.0 82.0 77.0 80.0wt % Corundum 12.0 9.2 15.0 1.9 4.4 1.6 wt % Rutile 0 0 0 1.0 1.0 traceWt % (Zr,Ti,Sn)O₄ 32.0 29.0 33.0 11.0 16.0 12.0 Wt % ZrO₂ 0 0 0 4.3 1.26.9 Pseudobrookite Lattice Parameter a 9.498 9.502 9.510 9.501 9.4779.471 Pseudobrookite Lattice Parameter b 9.724 9.729 9.738 9.729 9.7009.693 Pseudobrookite Lattice Parameter c 3.600 3.600 3.602 3.603 3.5973.597 (Zr,Ti,Sn)O₄ Lattice Parameter a 4.830 4.816 4.824 4.839 4.8434.856 (Zr,Ti,Sn)O₄ Lattice Parameter b 5.571 5.595 5.592 5.559 5.5425.503 (Zr,Ti,Sn)O₄ Lattice Parameter c 5.096 5.104 5.106 5.093 5.0855.073 Approximate Zr per ZrTiO₄ from unit cell 0.99 0.90 0.94 1.04 1.071.15 Approximate Ti per ZrTiO₄ from unit cell 0.61 0.58 0.56 0.61 0.640.67 Approximate Sn per ZrTiO₄ from unit cell 0.40 0.51 0.49 0.35 0.290.18 Al per Psb formula unit from EPMA 1.824 1.844 1.812 — — — Ti perPsb formula unit from EPMA 0.982 0.956 0.972 — — — Sn per Psb formulaunit from EPMA 0.064 0.080 0.076 — — — Zr per Psb formula unit from EPMA0.032 0.031 0.034 — — — Mg per Psb formula unit from EPMA 0.108 0.0980.118 — — — Zr per ZTT formula unit from EPMA 0.958 0.842 0.908 — — — Tiper ZTT formula unit from EPMA 0.619 0.576 0.572 — — — Sn per ZTTformula unit from EPMA 0.384 0.538 0.479 — — — Al per ZTT formula unitfrom EPMA 0.039 0.044 0.041 — — — Mg per ZTT formula unit from EPMA0.000 0.000 0.000 — — — Oxygen per ZTT formula unit from EPMA 3.9803.980

TABLE 6 Exemplary embodiments that contain a zirconium tin titanatephase for which MOR/S <1.0 Example Number 24 25 26 27 28 29 InorganicRaw Materials Al₂O₃ 12.0 μm  28.33 28.09 27.72 34.93 26.20 30.41 ZrO₂ A5.7 μm 0 0 0 0 0 0 ZrO₂ S 17.3 μm  24.98 20.97 24.44 10.03 20.07 17.20SnO₂ 5.5 μm 14.25 18.77 18.52 7.36 14.73 18.90 TiO₂ 0.5 μm 28.64 28.4025.60 40.58 33.69 30.59 Spinel 6.5 μm 3.81 3.78 3.73 7.10 5.32 2.90Firing Conditions Hold Temperature (° C.) 1400 1400 1400 1400 1400 1500Hold Time (hours) 15 15 15 15 15 15 Fired Properties % Porosity 65.064.5 63.5 59.9 61.8 56.1 d₁₀ 5.9 5.7 6.8 6.7 6.8 11.7 d₅₀ 11.7 11.8 11.813.0 12.4 14.4 d₉₀ 12.8 12.9 13.1 14.2 13.6 15.2 (d₅₀ − d₁₀)/d₅₀ 0.490.52 0.43 0.49 0.45 0.19 (d₉₀ − d₅₀)/d₅₀ 0.09 0.09 0.11 0.09 0.09 0.05(d₉₀ − d₁₀)/d₅₀ 0.59 0.61 0.54 0.58 0.54 0.24 CTE_(RT-1000° C.) (10⁻⁷°C.⁻¹) 46.3 40.3 48.7 18.0 29.5 17.9 CTE_(500-1000° C.) (10⁻⁷° C.⁻¹) 54.148.4 57.3 27.1 38.5 27.7 MOR (psi) 525 436 569 262 395 339 “S” value(psi) 577 541 616 420 490 408 MOR/S ratio 0.91 0.80 0.92 0.62 0.81 0.83Elastic Modulus at RT (psi) 5.37E+05 4.66E+05 6.32E+05 2.27E+05 3.38E+052.60E+05 MOR/(E*CTE_(500-1000° C.)) * = estimated 181 193 157 426 304470 Proportion and Composition of Phases Wt % Pseudobrookite 43.0 47.038.0 76.6 58.9 76.0 Wt % Corundum 20.0 18.0 22.0 6.6 10.0 6.2 Wt %Rutile 0.0 2.2 0.0 2.9 3.7 3.0 Wt % (Zr,Ti,Sn)O₄ 37.0 33.0 40.0 14 2715.0 Pseudobrookite Lattice Parameter a 9.500 9.499 9.509 9.484 9.4959.482 Pseudobrookite Lattice Parameter b 9.722 9.721 9.733 9.704 9.7199.707 Pseudobrookite Lattice Parameter c 3.605 3.603 3.605 3.602 3.6033.596 (Zr,Ti,Sn)O₄ Lattice Parameter a 4.801 4.788 4.799 4.773 4.7864.811 (Zr,Ti,Sn)O₄ Lattice Parameter b 5.560 5.584 5.582 5.541 5.5645.586 (Zr,Ti,Sn)O₄ Lattice Parameter c 5.076 5.085 5.089 5.052 5.0705.096 Approximate Zr per ZrTiO₄ from unit cell 0.88 0.80 0.85 0.78 0.810.89 Approximate Ti per ZrTiO₄ from unit cell 0.75 0.72 0.69 0.92 0.800.63 Approximate Sn per ZrTiO₄ from unit cell 0.37 0.48 0.47 0.30 0.390.48

Examples 14B, 17B, and 30

The stability of three aluminum titanate based ceramic bodies in thepresence of copper oxide was assessed at 1100° C. for inventive Examples14 and 17 and also for a ceramic comprised of magnesium-free aluminumtitanate+corundum+SrAl₂Si₂O₈. The results are provided in Table 7, whichshows the percentages of phases and their unit cell dimensions andvolumes before and after mixing the pulverized ceramic with 0.25 wt %copper (II) oxide and heating at 1100° C. for 2 hours. Themagnesium-free aluminum titanate+zirconium tin titanate ceramic ofExample 14B contained virtually no pseudobrookite-type phase after heattreatment with copper oxide, with the aluminum titanate havingdecomposed into corundum+rutile. By contrast, Example 17B, in which thepseudobrookite phase comprised approximately 92 mole % Al₂TiO₅ and 8mole % MgTi₂O₅, retained 92% of the initial pseudobrookite present inthe ceramic after testing, showing the benefit of magnesium on thestability of the inventive ceramics in the presence of copper oxide.Example 30, which relies on the presence of silicon to stabilize thepseudobrookite phase, retained only 32% of the initial aluminum titanatephase after heat treatment in the presence of copper.

TABLE 7 EX 14B EX 17B EX 30 Firing Conditions Hold Temperature (° C.)1500 1500 1410 Hold Time (hours) 15 15 16 As-Fired Properties Wt %Pseudobrookite (Psb) 59.8 69.2 72.8 Molar Ratio of Al₂TiO₅/MgTi₂O₅ in100/0 92/8 100/0 Psb Wt % Corundum 6.6 3.6 6.1 Wt % Rutile 1.0 1.0 0 Wt% (Zr,Ti,Sn)₂O₄ Solution 35.6 26.3 0 Wt % SrAl₂Si₂O₈ 0 0 20.7Pseudobrookite Lattice Parameter a 9.457 9.497 — Pseudobrookite LatticeParameter b 9.680 9.726 — Pseudobrookite Lattice Parameter c 3.591 3.601— Pseudobrookite Cell Volume (Å³) 328.7 332.6 — (Zr,Ti,Sn)₂O₄ LatticeParameter a 4.813 4.820 — (Zr,Ti,Sn)₂O₄ Lattice Parameter b 5.569 5.574— (Zr,Ti,Sn)₂O₄ Lattice Parameter c 5.087 5.094 — (Zr,Ti,Sn)₂O₄ CellVolume (Å³) 136.3 136.9 — After Copper Stability Test Wt %Pseudobrookite 0.3 63.8 23.4 Wt % Corundum 36.0 6.1 32.3 Wt % Rutile31.1 4.4 22.4 Wt % (Zr,Ti,Sn)₂O₄ Solution 32.8 25.8 0 Wt % SrAl₂Si₂O₈ 00 21.9 % of Initial Pseudobrookite Remaining 0.5% 92% 32% AT Lattice a(Å) — 9.501 — AT Lattice b (Å) — 9.729 — AT Lattice c (Å) — 3.602 — ATLattice Vol ((Å³) — 332.9 — ZTS Lattice Parameter a 4.760 4.822 — ZTSLattice Parameter b 5.538 5.574 — ZTS Lattice Parameter c 5.047 5.094 —ZTS Cell Volume (Å³) 133.1 136.9 —

What is claimed is:
 1. A method for preparing a ceramic body, saidmethod comprising the steps or: providing a batch composition comprisingat least one zirconium source, at least one tin source, at least onetitanium source, at least one aluminum source, and at least onemagnesium source; and firing the batch composition under conditionssuitable to form a ceramic body comprising at least one phase comprisinga pseudobrookite-type crystal structure and at least one phasecomprising zirconium tin titanate.
 2. The method according to claim 1,wherein the at least one magnesium source comprises MgAl₂O₄.
 3. Themethod according to claim 1, wherein the batch composition furthercomprises at least one pore forming agent.
 4. The method according toclaim 3, wherein the at least one pore forming agent comprises at leastone of graphite, starch, nut shells, and synthetic organic particles. 5.The method according to claim 4, wherein the starch comprises at leastone of sago palm starch, green mung bean starch, canna starch, cornstarch, rice starch, pea starch, and potato starch.
 6. The methodaccording to claim 3, wherein the at least one pore forming agentcomprises a median particle diameter in a range of 1 to 60 microns. 7.The method according to claim 1, wherein the batch is fired at atemperature of at least about 1400° C.
 8. The method according to claim1, further comprising extruding the batch composition to form ahoneycomb structure.
 9. The method according to claim 1, furthercomprising disposing a catalyst on the ceramic body.
 10. The methodaccording to claim 1, wherein the ceramic body comprises about 30 wt %to about 90 wt % of the at least one phase comprising apseudobrookite-type crystal structure and about 10 wt % to about 45 wt %of the at least one phase comprising zirconium tin titanate.
 11. Themethod according to claim 1, wherein the at least onepseudo-brookite-type crystal structure comprises magnesium-containingAl₂TiO₅.
 12. The method according to claim 1, wherein the at least onephase comprising zirconium tin titanate is of the formula(Al)_(p1)(Zr)_(q1)(Ti)_(q2)(Sn)_(q3)O_(4-0.5(p1)) wherein p1+q1+q2+q3=2;p1 ranges from about 0 to about 0.08; the value of q1/(q1+q2+q3) rangesfrom about 0.36 to about 0.60; the value of q2/(q1+q2+q3) ranges fromabout 0.23 to about 0.50; and the value of q3/(q1+q2+q3) ranges fromabout 0.05 to about 0.33.
 13. The method according to claim 1, whereinthe ceramic body comprises less than about 15% of a corundum phase. 14.The method according to claim 1, wherein the ceramic body comprises aporosity of at least about 45%.